High strength spring-use heat treated steel

ABSTRACT

The present invention provides spring-use heat treated steel which is cold coiled, can achieve both sufficient atmospheric strength and coiling workability, has a tensile strength of 2000 MPa or more, and can improve the performance as a spring by heat treatment after shaping the spring, that is, heat treated steel for high strength spring containing, by mass %, C: 0.45 to 0.9%, Si: 1.0 to 3.0%, Mn: 0.1 to 2.0%, V: over 0.1 to 1.0%, N: limited to 0.007% or less, and a balance of Fe and unavoidable impurities, and satisfying, in terms of the analyzed value of the extracted residue after heat treatment, [amount of V in filtrate filtered by 0.2 μm filter (mass %)]≧[amount of V in steel (mass %)]×0.4.

TECHNICAL FIELD

The present invention relates to hot or cold coiled, high strength and high toughness, in particular cold coiled and nitrided high strength heat treated spring steel.

BACKGROUND ART

Along with the reduction in weight and improvement in performance of automobiles, springs are being made higher in strength. High strength steel having a tensile strength exceeding 1500 MPa after heat treatment is being used for springs. In recent years, steel wire having a tensile strength exceeding 2100 MPa is also being sought. This is so as to secure a hardness of material where even with some softening due to stress-relief annealing, nitridation, and other heating at the time of spring production, there is no problem for the spring.

Further, it is known that with nitridation or shot peening, the surface hardness rises and the durability during spring fatigue is remarkably improved, but the spring setting characteristic is not determined by the surface hardness. The internal strength or hardness of the spring material also has a great effect. Therefore, it is important to design ingredients able to maintain the internal hardness extremely high.

As a technique for this, there is an invention adding V, Nb, Mo, or another element to form fine carbides dissolving by quenching and precipitating by tempering and thereby limiting the movement of dislocations and improving the anti-setting property (for example, see Japanese Patent Publication (A) No. 57-32353).

On the other hand, among the methods for production of steel coil springs, there are “hot coiling” comprising heating the steel to the austenite region for coiling, then quenching and tempering it and “cold coiling” comprising quenching and tempering the steel in advance and cold coiling the resultant high strength steel wire. With cold coiling, it is possible to use oil tempering, high frequency treatment, etc. enabling rapid heating and rapid cooling at the time of production of the steel wire, so it is possible to reduce the grain size of the prior austenite of the spring material. As a result, it is possible to produce a spring superior in fracture characteristics. Further, it is possible to simplify the heating furnace and other facilities on the spring production line, so there is the advantage to the spring manufacturers as well that this leads to a reduction in the capital costs. Recently, springs are being cold worked. In suspension springs, compared with valve springs, thicker steel wires are used, but cold coiling is introduced due to its advantages.

However, if the spring-use steel wire for cold coiling increases in strength, it will break at the time of cold coiling and will be unable to be formed into a spring shape in many cases. Up to now, both strength and workability could not be achieved, so the wire had to be coiled by industrially disadvantageous methods of hot coiling or coiling, then quenching and tempering to obtain both strength and workability.

Further, when cold coiling and nitriding high strength heat treated steel wire to secure its strength, it has been believed effective to add large amounts of V, Nb, and other so-called alloy elements causing precipitation of fine carbides in the steel. However, if actually adding large amounts, they will not dissolve by the heating at the time of quenching, but will grow coarser and form so-called undissolved carbides which will become factors behind fracture at the time of cold coiling. Therefore, technology focusing on the undissolved carbides has also been seen.

There is an invention aiming at improvement of performance by controlling not only the alloy elements, but also the carbides such as the cementite present in large amounts in the steel (for example, see Japanese Patent Publication (A) No. 2002-180198).

These patents prescribe the spherical carbides in detail and try to achieve both workability and higher spring strength, but even if suppressing the carbides of such relatively clear spherical carbides (alloy-based and cementite-based), there are limits to further increase of the strength and improvement of the spring performance. That is, these provisions by nature strongly suppress “defects” and suppress deterioration of the workability. There were also limits to the direct strengthening of the spring performance.

DISCLOSURE OF THE INVENTION

The present invention has as its object the provision of spring use heat treated steel which is cold coiled, can achieve both sufficient ordinary temperature strength and coilability, has a tensile strength of 2000 MPa or more, and can improve the performance as a spring by heat treatment after spring fabrication.

The inventors controlled the N, which had never been taken note of up to now, to suppress the formation of undissolved carbides even with the addition of alloy elements and achieve toughness and, as workability, high strength and coilability at the same time, and used subsequent nitridation and other heat treatment to utilize to the fullest the precipitation type fine carbides never before considered so as to develop heat treated steel able to improve the spring performance.

That is, the present invention has as its gist the following:

-   -   (1) Spring steel characterized by containing, by mass %,         -   C, 0.45 to 0.9%,         -   Si: 1.0 to 3.0%,         -   Mn: 0.1 to 2.0%,         -   V: over 0.1 to 1.0%,         -   N: limited to 0.007% or less, and         -   a balance of Fe and unavoidable impurities, and         -   satisfying, in terms of the analyzed value of the extracted             residue after heat treatment,

[amount of V in filtrate filtered by 0.2 μm filter(mass %)]≧[amount of V in steel(mass %)]×0.4.

-   -   (2) Spring-use heat treated steel wire as set forth in (1)         characterized by further containing, by mass %, one or more of         Cr: 0.5 to 2.5%, Nb: 0.001 to less than 0.05%, Ti: 0.001 to less         than 0.05%, W: 0.05 to 0.5%, Mo: 0.05 to 0.5%, Ta: 0.001 to         0.5%, Ni: 0.05 to 3.0%, Cu: 0.05 to 0.5%, Co: 0.05 to 3.0%, B:         0.0005 to 0.006%, Te: 0.0002 to 0.01%, Sb: 0.0002 to 0.01%, Mg:         0.0001 to 0.0005%, Zr: 0.0001 to 0.0005%, Ca: 0.0002 to 0.01%,         and Hf: 0.0002 to 0.01%.     -   (3) Spring-use heat treated steel wire as set forth in (1)         or (2) characterized further by restricting, by mass %, Al to         0.005% or less.     -   (4) Spring-use heat treated steel characterized by having the         steel ingredients as set forth in any of (1) to (3), having a         prior austenite grain size number after quenching and tempering         of #10 or more, and having residual austenite in an amount of 15         mass % or less.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a schematic view for explaining a filtrate passing through a 0.2 μm filter in analysis of V by electrolysis (speed method).

FIG. 2 is a view showing the relationship between the annealing temperature and tensile characteristics (tensile strength and yield point) for showing the effect of the “amount of V (mass %) in filtrate filtered by a 0.2 μm filter”.

BEST MODE FOR CARRYING OUT THE INVENTION

The inventors prescribed the chemical compositions for obtaining high strength and further heat treated the steel to control the shapes of the carbides in the steel and secure sufficient coilability for production of springs and thereby invented a spring-use heat treated steel enabling improvement of the spring performance by annealing or other heat treatment after being spring fabrication. The details are shown below. First, the chemical compositions of the steel will be explained.

C: 0.4 to 0.9%

C is an element having a large effect on the basic strength of the steel material. To obtain a strength more sufficient than the past, the amount was made 0.4 to 0.9%. If less than 0.4%, a sufficient strength cannot be obtained. In particular even in the case of eliminating the nitridation for improving the spring performance, to secure a sufficient spring strength, 0.4% or more of C is preferable. If over 0.9%, de facto over coprecipitation results and a large amount of coarse cementite precipitates, so the toughness is remarkably reduced. This simultaneously reduces the coiling characteristics. Therefore, the upper limit of the amount of C was made 0.9%.

Further, the relationship with the microstructure is also close. If less than 0.4%, the number of carbides is small, so the area ratio of regions where the carbide distribution is locally smaller than other parts (hereinafter referred to as “carbide-poor regions”) easily increases and a sufficient strength and toughness or coilability (ductility) are difficult to obtain. Therefore, the amount is preferably made 0.55% or more, from the viewpoint of the balance of strength and coiling, is more preferably made 0.6% or more.

On the other hand, when the amount of C is large, dissolution of alloy-based or cementite-based carbides by the heating at the time of quenching tends to become difficult. When the heating temperature at the time of heat treatment is high and when the heating time is short, the strength and the coilability are often insufficient.

Further, the undissolved carbides also have an effect on the carbide-poor regions. If C forms undissolved carbides in the steel, the amount of substantive C in the matrix falls and the area ratio of so-called carbide-poor regions where the distribution of carbides becomes smaller than other locations sometimes increases. These carbide-poor regions reduce the mechanical properties, so must be avoided at all costs. For this reason as well, undissolved carbides and other uneven distribution of C in the steel are preferably avoided.

Further, if the amount of C is increased, in the case of medium carbon steel, the form of the martensite at the time of tempering becomes general lath martensite, while if the amount of C is large, it is known that the form changes to lenticular martensite. As a result of research and development, the inventors discovered that the carbide distribution of a tempered martensite structure formed by tempering lenticular martensite is lower in carbide density compared with one obtained by tempering lath martensite. Therefore, by increasing the amount of C, the increase in the lenticular martensite or undissolved carbides causes the carbide-poor regions to increase in some cases.

For this reason, the content is preferably made 0.7% or less, more preferably 0.65% or less, whereby the undissolveds carbide or carbide-poor regions can be reduced relatively easily.

Si: 1.0 to 3.0%

Si is an element necessary for securing the spring strength and the hardness and anti-setting characteristic. If less than this, the necessary strength and anti-setting characteristic are insufficient, so 1.0% was made the lower limit. Further, Si has the effect of making the Fe carbide-based precipitates at the grain boundaries spherical and finer. There is an effect of making the Fe carbides finer and reducing the occupied area ratio of the grain boundary precipitates at the grain boundaries. However, if added in too large an amount, not only is the material made to harden, but also it is made brittle. Therefore, to prevent embrittlement after quenching and tempering, 3.0% was made the upper limit. Here, “iron-based carbides” include, in addition to the so-called cementite, the Fe₂₋₃C known as ε-carbide etc.

Si is also an element contributing to tempering softening resistance, so to produce a high strength wire material, a large amount is preferably added to a certain extent. Specifically, adding 2% or more is preferable. On the other hand, to obtain a stable coilability, preferably 2.6% or less is added.

Mn: 0.1 to 2.0%

Mn is frequently used for deoxidation or immobilization of the S in the steel as MnS and improves the quenching ability to obtain sufficient hardness after heat treatment. To secure this stability, 0.1% is made the lower limit. Further, to prevent embrittlement by Mn, the upper limit was made 2.0%. Further, to achieve both strength and coilability, 0.3 to 1% is preferable. When giving priority to the coiling, making the amount 1.0% or less is effective.

V: over 0.1 to 1.0%

In the present invention, control of the V is a major point. Here, the amount of addition of V will be explained. V causes precipitation of carbides and hardening at the time of tempering for secondary precipitation hardening etc., so can be utilized for hardening the steel wire at the tempering temperature or for hardening the surface layer at the time of nitridation. Further, the formation of nitrides, carbides, and carbonitrides is effective for suppression of coarsening of the grain size of the austenite, so addition is preferable.

However, nitrides, carbides, and carbonitrides of V are also formed at the austenite transformation temperature A₃ point of steel or more, so when insufficiently dissolved, easily remain as undissolved carbides (nitrides). The undissolved carbides not only become a cause of fracture at the time of spring coiling, but also “wastefully consume V”. They reduce the effect of improvement of the tempering softening resistance and secondary precipitation hardening due to the added V and end up reducing the spring performance. Therefore, up until now, the preferable amount of addition of V was considered to be 0.1% or less industrially.

However, in the present invention, the amount of N is controlled to suppress the formation of V-based nitrides, carbides, and carbonitrides at the austenite transformation temperature A₃ point or more, so that much more V can be added. The amount of V added was therefore made over 0.1% to 1.0%. If the amount of addition is 0.1% or less, there is little effect of improvement of the hardness of the nitride layer, increase of the depth of the nitride layer, or other effects of addition of V, so the content is made over 0.1%, more preferably 0.15% or more. Further, if the amount of addition is over 1.0%, coarse undissolved inclusions are formed, the toughness is reduced, and, in the same way as Mo, a supercooled structure easily results which easily causes cracking or fracture at the time of drawing. Therefore, 1.0%, where industrially stable handling becomes easy, was made the upper limit.

Nitrides, carbides, and carbonitrides of V are also formed at the austenite transformation temperature A₃ point or more of steel, so when these are insufficiently dissolved, they easily remain as undissolved carbides (nitrides). Therefore, if considering the current industrial ability to control the nitrogen content, industrially the content is preferably made 0.5% or less, more preferably 0.4% or less.

On the other hand, in surface hardening by nitridation, the material is reheated to a temperature of 300° C. or more, so to suppress the hardening of the surfacemost layer or softening of the internal hardness due to nitridation, addition of over 0.1% is necessary. 0.15% or more is preferable, further 0.2% or more added is preferable. These points will be explained in further detail in the section on the explanation of the designation of the amount of V on the filter.

N: Restricted to 0.007% or Less

In the present invention, a strict limit value of N≦0.007% is prescribed. The effects of N in steel are as follows: (1) It is present as dissolved N in ferrite and suppresses movement of dislocations in the ferrite to thereby make the ferrite harden, (2) it forms nitrides with Ti, Nb, V, Al, B, and other alloy elements to affect the performance of the steel material. The mechanism etc. will be explained later. (3) It affects the precipitation behavior of cementite and other iron-based carbides and affects the performance of the steel material.

In spring steel, the strength is secured by alloy elements such as C, Si, and V, so the effect of hardening by the dissolved N is not large. On the other hand, when considering the cold working (coiling) of the spring, it suppresses movement of dislocations and thereby suppresses deformation of the worked parts and causes embrittlement of the worked parts, so lowers the coiling characteristic.

Further, among the elements defined in claim 1, V forms precipitates in the steel at a high temperature. The chemical ingredients are mainly nitrides at a high temperature. Along with cooling, they are changed in form to carbonitrides and carbides. Therefore, the nitrides formed at a high temperature easily become nuclei for the precipitation of V carbides. This means that undissolved carbides are easily formed at the time of heating in the patenting or quenching process. These form nuclei, so easily grow in size.

Further, if viewed from the viewpoint of cementite, in the high strength spring such as the present invention, from the required strength, the tempering is performed at a tempering temperature of 300 to 500° C. In spring steel, from the characteristic ingredients, the iron-based carbides formed at the time of tempering are made to change complicatedly in form to ε-carbide or θ-carbide (so-called cementite Fe₃C). Therefore, this has an effect on the ductility and other mechanical properties of the steel. N also has an effect on the formation of its carbides. If the amount of N is small, the ductility and toughness at 350 to 500° C. are improved.

Further, if N exceeds 0.007%, V-based nitrides are easily formed, a large amount of undissolved carbides are formed, and, depending on the form of the ferrite or carbides, the steel is embrittled.

In the present invention, to reduce the damage by such N, the amount of N is limited to 0.007% or less. Further, the amount of N is preferably kept down to 0.004% or less. Further, as explained later, it is also effective to add slight amounts of one or more of Ti, Ta, and Nb.

Originally, if suppressing the amount of N to 0.004% or less, good performance can sometimes be obtained even without adding one or more of Ti, Ta, and Nb, but sometimes it is difficult to industrially stably reduce the amount to 0.004% or less or this is disadvantageous in terms of the production costs. Therefore, addition of a slight amount of one or more of Ti, Ta, and Nb proves to be a realistic technique.

If adding a slight amount of one or more of Ti, Ta, and Nb, since these elements form nitrides at a high temperature, the dissolved nitrogen is de facto reduced, so a similar effect can be obtained as with reducing the amount of addition of N. Therefore, it is also possible to raise the upper limit of the amount of addition of N. However, if the amount of N exceeds 0.007%, the amount of formation of V, Nb, or Ti nitrides would become too great. As a result, the amount of undissolved carbides would become greater and TiN and other hard inclusions would increase, so the toughness would fall and the fatigue durability characteristics and coiling characteristics would fall. Therefore, the upper limit of the amount of N was restricted to 0.007%.

That is, even when adding one or more of Ti, Ta, and Nb, if the amount of N becomes too great or the amount of addition of one or more of Ti, Ta, and Nb becomes too great, coarse nitrides of Ti, Ta, and Nb would again be formed and conversely would become harmful, so the amount of addition of the one or more of Ti, Ta, and Nb has to be kept slight.

For this reason, the upper limit of the amount of N is preferably 0.005% or less, more preferably 0.004% or less.

This precision control of the N suppresses ferrite embrittlement and suppresses formation of V-based nitrides and thereby suppresses the formation and growth of undissolved carbides. Further, controlling the form of the iron-based carbides enables improvement of the toughness.

In this way, even when adding Ti, Ta, or Nb, if considering the ease of heat treatment etc., the amount of N is preferably made 0.005% or less. The smaller the amount of N the better. While substantially 0% is also all right, N easily is mixed in from the atmosphere in the steelmaking process etc., so considering the production costs and the ease in the denitridation process, the amount is preferably 0.0015% or more.

The provision that, in terms of the analyzed value of the extracted residue after heat treatment, [amount of V in filtrate filtered by 0.2 μm filter (mass %)]≧[amount of V in steel (mass %)]×0.4 will be explained next.

This provision is a key point of the present invention. In the past, the clearly observable spherical undissolved carbides had mainly been prescribed, but this provision not only does that, but also enables a very large amount of the precipitated V carbides effective for the final spring performance to be obtained by suppressing the coarse V-based carbides, regardless of the spherical or aspherical shape or dissolved or undissolved precipitation history, and thereby promoting the dispersion of a large amount of fine V-based carbides after spring fabrication and improving spring performance through control of the steel.

A small austenite grain size is better for fatigue durability, so reducing the heating temperature at the time of quenching is effective. However, this simultaneously increases the undissolved carbides, so it is necessary to form fine precipitates or dissolved V effective even after annealing like in the present invention. For this reason, attention must also be paid to the heat treatment conditions of the processes up to quenching such as rolling and patenting so as to enable sufficient dissolution even at heating temperatures of relative low temperatures where the austenite grain size can be reduced.

Considering the presence of such undissolved carbides, the inventors discovered that if filtering a solution electrolyzed at a constant potential by a filter, the V-based carbides remaining on the filter were undissolved carbides or similar coarse carbides and that the V improving the properties of the steel after quenching and tempering very probably passed through the filter.

Specifically, the inventors discovered that if the [amount of V in filtrate filtered by 0.2 μm filter (mass %)] after quenching and tempering is greater than [amount of V in steel (mass %)]×0.4, both strength and workability can be achieved before annealing without impairing the spring workability and that, further, the amount of dissolved V or fine V-based precipitates precipitating at the time of heating and annealing of the nitridation etc. is great, the mechanical properties (yield point) after annealing are great, and the spring performance is superior.

Here, the method of measuring the [amount of V in filtrate filtered by 0.2 μm filter (mass %)] will be explained. FIG. 1 is a schematic view for explaining the filtrate passing through a 0.2 μm filter in analysis of V by electrolysis (speed method). As shown in FIG. 1, in the present invention, quenched and tempered heat treated steel wire is electrolyzed to dissolve the ferrite component and obtain an electrolytic solution 1 by the speed method. This solution is filtered 2 to obtain an extract residue 4 on a filter 3. Further, V is also present in the filtrate 5 filtered by the 0.2 μm filter.

For the electrolysis, the so-called speed method is used. This method is also used for preparation of replica samples for transmission type electron microscopes for observing ferrous materials. The potential and solution are strictly controlled during the electrolysis so as to enable the ferrite component to be preferentially electrolyzed. Specifically, a constant potential electrolysis system using a Model FV-138 electrolysis apparatus made by Fujiwara Seisakusho is used. The solution is a commercially available speed method electrolytic solution (product name Electrolyte A). After 2000 coulomb electrolysis, the obtained electrolytic solution is suction filtered by a 0.2 μm mesh filter to obtain the residue. Further, it is possible to precisely measure the mass of the steel wire before and after electrolysis to measure the [electrolysis amount (mass)]. By measuring the amount of V in the residue on the filter and dividing it by the electrolysis amount, it is possible to obtain a grasp of the relative amount of coarse V-based carbides (mass %) remaining on the filter among the added amount of V.

The amount of V in the residue on the filter can be measured by spectrometry (ICP) based on Attachment 1 of JIS G 1258-1999. [Amount of V in residue on filter (mass)]/[electrolysis amount (mass)]×100=[amount of V-based carbides remaining on filter (mass %)], so the value of this subtracted from the amount of V added in the steel (mass %) was made the [amount of V in filtrate filtered by 0.2 μm filter (mass %)].

$\begin{matrix} {{{{Amount}\mspace{14mu} {of}\mspace{14mu} V\mspace{14mu} {in}\mspace{14mu} {steel}\mspace{20mu} \left( {{mass}\mspace{14mu} \%} \right)} - {\frac{\begin{matrix} {{Amount}\mspace{14mu} {of}\mspace{14mu} V} \\ {{in}\mspace{14mu} {filter}\mspace{14mu} {residue}} \end{matrix}}{{Steel}\mspace{14mu} {electrolysis}} \times 100}} = {\frac{\begin{matrix} {{Filtrate}\mspace{14mu} V} \\ {amount} \end{matrix}}{{Steel}\mspace{11mu} {electrolysis}} \times 100}} \\ {= \begin{matrix} \begin{matrix} \begin{matrix} {{Amount}\mspace{14mu} {of}\mspace{14mu} V} \\ {{in}\mspace{14mu} {filtrate}\mspace{14mu} {passing}} \end{matrix} \\ {0.2\mspace{14mu} {µm}\mspace{14mu} {filter}} \end{matrix} \\ \left( {{mass}\mspace{14mu} \%} \right) \end{matrix}} \end{matrix}$

In general, the smaller the mesh, the more coarse precipitates remain on the filter as residue, but in practice, even with a filter of 0.2 μm mesh, precipitates of 0.2 μm or less remain and are obtained. Due to this, in the present invention, even 0.2 μm or less precipitates are taken note of for finer carbide control. Therefore, in analysis of the solution passing through a 0.2 μm filter, the amount of V in the dissolved state or contained in extremely fine precipitates smaller than the filter mesh size is measured.

V passing through such a 0.2 μm filter raises the strength of the treated steel wire, but still will not degrade the workability at the time of spring fabrication. Further, even in the nitridation after spring fabrication, the dissolved V will form extremely fine precipitates, so despite being heated to the extent of nitridation, the yield point of the material is raised. When that effect is large, the tensile strength and hardness will sometimes also be raised. In particular, if reheating steel to 450 to 650° C. after quenching, it is possible to cause the precipitation of a large amount of very fine V-based carbides able to pass through a 0.2 μm mesh. The heating of the nitridation process is also utilized, so by making the amount of V inside the spring the prescribed amount, it is possible to improve the spring performance. This is important not only for just suppressing factors causing deterioration such as the prevention of the deterioration of workability like in the inventions up to now, but also in positively improving the spring performance.

Cr: 0.5 to 2.5%

Cr is an effective element for improving the quenching ability and the tempering softening resistance. Further, in the nitridation such as seen in recent high strength valve springs, it is an element effective not only for securing tempering hardness, but also increasing the surface layer hardness after nitridation and the depth of the hardened layer. However, if the amount of addition is large, not only is an increase in cost incurred, but also the cementite seen after the quenching and tempering is made coarser. Further, the alloy-based carbides are stabilized and made coarser as an effect. As a result, the wire material becomes brittle, so easily breaks at the time of coiling. Therefore, when adding Cr, if not 0.5% or more, the effect is not clear. Further, 2.5%, where the embrittlement becomes remarkable, was made the upper limit.

However, in the present invention, by prescribing the N, the carbides are controlled to become fine, so a large amount of Cr can be added. The amount of addition easily giving a high strength was therefore used.

Further, when performing nitridation, addition of Cr enables the nitridation hardened layer to be made deeper. Therefore, addition of 1.1% or more is preferable. Further, addition of 1.2% or more is preferable for nitridation for high strength springs never before seen.

Cr obstructs the melting of the cementite by heating, so if the amount of C becomes a large C>0.55%, suppressing the amount of Cr enables the formation of coarse carbides to be suppressed and both strength and coilability to be easily achieved. Therefore, preferably the amount of addition is made 1.7% or less.

Nb: 0.001 to less than 0.05%

Nb forms nitrides, carbides, and carbonitrides. The nitrides are formed at a higher temperature than V. For this reason, by forming Nb nitrides at the time of cooling to bond with the N in the steel, the V-based nitride formation temperature is lowered. Therefore, in the numerous heat treatment processes applied to a material up to the spring fabrication, it is possible to suppress coarsening of the V-based carbonitrides and promote dissolution in the heating process of the transformation point or higher. As a result, the formation of V-based undissolved carbides can be suppressed, so the spring workability of the high strength steel wire and the tempering softening resistance by the V-based precipitates after spring fabrication can be effectively secured.

Further, in addition to the suppression of coarsening of the grain size of austenite by Nb-based carbonitrides, these may be used for hardening the steel wire in the tempering temperature and for hardening the surface layer at the time of nitridation. However, if the amount of addition is too great, undissolved carbides having Nb-based nitrides as nuclei will easily remain, so addition of a large amount should be avoided. Specifically, if the amount of addition of Nb is less than 0.001%, almost no effect of addition can be observed. Further, if 0.05% or more, addition of a large amount would form coarse undissolved inclusions and lower the toughness and, like with Mo, would easily cause the formation of a supercooled structure which would easily cause cracking or fracture during drawing. Therefore, the amount was made less than the 0.05% where industrially stable handling is easy.

Further, Nb inherently lowers the hot ductility and easily causes defects in the rolling process, so addition of the minimum required amount is preferable. The amount of addition is preferably 0.03% or less, more preferably 0.015% or less.

Ti: 0.001 to less than 0.05%

In the present invention, when adding Ti, the amount of addition should be 0.001% to less than 0.05%. Ti is a deoxidizing element and a nitride- and sulfide-forming element, so has an effect on the formation of oxides, nitrides, and sulfides. Therefore, a large amount of addition would easily cause the formation of hard oxides and nitrides, so if added incautiously, it would form hard carbides and lower the fatigue durability. In the same way as Al, in particular in high strength springs, reduces the variation of the fatigue strength rather than the fatigue limit of springs itself. If the amount of Ti is large, the rate of fracture due to inclusions becomes greater, so the amount has to be controlled and was made less than 0.05%.

On the other hand, Ti forms TiN in the molten steel at a high temperature, so acts to reduce the sol.N in the molten steel. In the present invention, the technical point is to restrict the N so as to suppress the formation of V-based nitrides and suppress the growth of V-based undissolved carbides. For this reason, if consuming the N at a temperature of the V-based nitride formation temperature or more, it is possible to suppress the growth of V-based nitrides and V-based carbonitrides growing at the time of cooling using these as nuclei. That is, by adding Ti, the amount of N de facto bonding with the V is reduced, so the temperature of formation of V-based nitrides is lowered and further the V-based undissolved carbides can be suppressed.

Therefore, addition of a large amount of Ti should be avoided from the viewpoint of formation of Ti-based undissolved carbonitrides and oxides, but addition of a slight amount can lower the V-based nitride formation temperature, so rather can reduce the undissolved carbides. The amount of addition is 0.001% or more. If less than 0.001%, there is no effect of consumption of N, there is no effect of suppression of V-based undissolved carbides, and no effect of improvement of the workability is seen. However, the amount of addition of Ti is preferably 0.02% or less, more preferably 0.01% or less.

W: 0.05 to 0.5%

W improves the quenchability and acts to form carbides in steel to raise the strength, so is effective for imparting tempering softening resistance. Therefore, addition is preferable as much as possible.

W forms carbides at a lower temperature compared with Ti, Nb, etc., so does not easily form undissolved carbides, but can impart tempering softening resistance by precipitation hardening. That is, it hardly remains as undissolved carbides causing harm even in the heat treatment up to the spring fabrication. On the other hand, it never causes a large drop in internal hardness even in nitridation or strain relief annealing performed at a relatively low temperature.

If the amount of addition is less than 0.05%, no effect is seen. If over 0.5%, a supercooled structure easily forms or, when industrial heat treatment is performed, conversely the ductility or other mechanical properties are liable to be impaired, so the amount of addition of W was made 0.05 to 0.5%. Further, if considering the ease of heat treatment etc., 0.1 to 0.4% is preferable. In particular, to avoid a supercooled structure right after rolling or other trouble and obtain the maximum limit of tempering softening resistance, addition of 0.15% or more is more preferable.

Mo: 0.05 to 0.5%

Mo improves the quenchability and precipitates as carbides at a heat treatment temperature of a relatively low temperature of the extent of tempering and nitridation temperature, so can easily give tempering softening resistance. Therefore, it is possible to give a high strength without softening even after high temperature tempering or strain relief annealing, nitridation, or other heat treatment performed in accordance with need before spring fabrication. This suppresses the drop in spring internal hardness after nitridation, so increases the effect of hot setting or strain relief annealing and improves the fatigue characteristics of the final spring. Specifically, it is possible to make the tempering temperature at the time of controlling the strength a high temperature. This making the tempering temperature a high temperature has the effect of making the grain boundary carbides precipitating in a film shape spherical by tempering at a high temperature and reducing the grain boundary area ratio. This is advantageous for securing the grain boundary strength and improving the delayed fracture and embrittlement fracture characteristics.

Mo forms Mo-based carbides separate from cementite in the steel. In particular, compared with V etc., the precipitation temperature is low, so there is the effect of suppressing coarsening of the carbides. If the amount of addition is less than 0.05%, the effect cannot be recognized. However, if the amount of addition is large, it is preferable to patent the steel material in advance at the time of drawing so as to make it a ferrite-pearlite structure. Despite this, a supercooled structure easily occurs due to the softening heat treatment before rolling or drawing etc. and cracking or fracture at the time of drawing is easily caused. This is because Mo is an element greatly imparting quenchability, so if the amount of addition becomes large, the time until the end of the pearlite transformation becomes long, the temperature cannot be maintained until the end of transformation at the time of cooling after rolling or in the patenting process, a supercooled structure easily occurs and causes fracture at the time of drawing, or when not breaking and remaining as internal cracks, the characteristics of the final product are greatly degraded. From this, if Mo is over 0.5%, the quenchability increases and industrially obtaining a ferrite-pearlite structure becomes difficult, so this is made the upper limit. To suppress the formation of a martensite structure reducing the producibility in the rolling, drawing, or other manufacturing process and facilitate industrially stable roll and drawing, the content is preferably made 0.4% or less, more preferably 0.2% or so.

Compared with V, Nb, and Ti having the effects of strengthening the tempering softening resistance in the same way as W and Mo, V, Nb, and Ti, as explained above, form nitrides and further cause easy growth of carbides using these as nuclei, while W and Mo do not form nitrides much at all, so can be added and strengthen the softening resistance without being affected by the amount of N. That is, V, Nb, and Ti can also strengthen the softening resistance, but if added for strengthening the softening resistance while avoiding undissolved carbides, the amount of addition ends up being inherently restricted. Therefore, when not forming undissolved carbides and requiring higher softening resistance, addition of W or Mo, which do not form nitrides, precipitate as carbides at a relatively low temperature, and function as precipitation strengthening elements, is extremely effective.

Ta: 0.001 to 0.5%

Ta forms nitrides, carbides, and composite precipitates and is effective for imparting tempering softening resistance and securing toughness by reducing the γ-grain size. If the amount of addition is less than 0.001%, the effect is not clear, while if over 0.5%, coarse nitrides, carbides, and their composite precipitates are formed and impair the workability etc., so this was made the upper limit.

Further, one or more of Ni, Cu, Co, and B is added to secure strength by matrix hardening when trying to achieve both strength and workability and the optimum balance of the softening resistance and workability due to carbide control cannot be obtained.

Ni: 0.05 to 3.0%

Ni does not form carbides or other precipitates, but can improve the quenching ability and stably increase the strength by heat treatment. Further, it can improve the ductility of the matrix and improve the coilability. However, with quenching and tempering, it increases the residual austenite, so the material is inferior in terms of setting of spring formation or uniformity of the material. If the amount of addition is 0.05% or less, no effect can be recognized in increasing the strength and improving the ductility. On the other hand, addition of a large amount of Ni is not preferable. At 3.0% or more, problems such as an increase in residual austenite becomes remarkable, the effect of improvement of the quenching ability and improvement of the ductility become saturated, and there are cost disadvantages etc. Residual austenite changes to martensite by work induced transformation when deformation is introduced and breaks down even with relatively low temperature treatment, so a stable material is difficult to maintain, so the less the better. For this reason, it is not preferable to excessively add NI—which easily causes residual austenite to remain.

Cu: 0.05 to 0.5%

Addition of Cu is effective for preventing the decarburization lowering the spring fatigue life after spring fabrication. Further, like Ni, it has the effect of improving the corrosion resistance. Usually, to stabilize the wire diameter and remove the decarburized layer, the surface layer is removed by peeling. By suppressing the decarburized layer, it is possible to improve the fatigue life of the spring and eliminate the peeling step. The effect of Cu in suppressing decarburization and the effect in improving the corrosion resistance can be exhibited when 0.5% or more. As explained later, even if adding Ni, if over 0.5%, embrittlement easily causes rolling defects. Therefore, the lower limit was made 0.05% and the upper limit was made 0.5%. The addition of Cu does not detract much at all from the mechanical properties at room temperature, but even if adding Cu over 0.3%, the hot ductility is degraded, so sometimes the billet surface cracks during rolling. For this reason, the amount of addition of Ni for preventing cracking during rolling is preferably made [Cu %]<[Ni %] in accordance with the amount of addition of Cu. In the range of Cu of 0.3% or less, rolling defects are not caused, so it is not necessary to limit the amount of addition of Ni for the purpose of preventing rolling defects.

Co: 0.05 to 3.0%

Co reduces the quenching ability in some cases, but improves the high temperature strength. Further, to inhibit the growth of carbides, it acts to suppress the formation of coarse carbides which become a problem in the present invention. Therefore, it is possible to suppress the coarsening of the carbides including cementite. Therefore, addition is preferable. When added, if 0.05% or less, the effect is small. However, if added in a large amount, the ferrite phase increases in hardness and reduces the ductility, so the upper limit was made 3.0%. This industrially gives a stable performance at 0.5% or less.

B: 0.0005 to 0.006%

B is an element improving the quenching ability and has an effect of cleaning the austenite grain boundaries. The addition of B renders harmless the P, S, and other elements segregating at the grain boundary and reducing the toughness and therefore improves the fracture characteristics. At this time, the effect is lost if the B bonds with N and forms BN. The lower limit of the amount of addition is made 0.0005% where the effect becomes clear, while the upper limit is made 0.0060% where the effect becomes saturated. However, if even a little amount of BN is formed, it causes embrittlement, so full consideration is required so as to not produce BN. Therefore, preferably the amount is 0.003 or less. More preferably, it is effective to immobilize the free N by the Ti or other nitride-forming elements and make the amount of B 0.0010 to 0.0020%.

These Ni, Cu, Co, and B are mainly effective for strengthening the ferrite phase of the matrix. They are elements effective for securing strength by matrix strengthening when trying to achieve both strength and workability simultaneously and the optimum balance of softening resistance and workability cannot be obtained by carbide control.

Al: restricted to 0.005% or less

Al is a deoxidizing element and has an effect on formation of oxides. In particular, in high strength valve springs, Al₂O₃ and other hard oxides easily form fracture starting points, so it is necessary to avoid this. Therefore, it is important to strictly control the amount of Al. In particular, when the tensile strength of the heat treated steel wire would exceed 2100 MPa, strict control of the oxide-forming elements is essential for reducing the variation in fatigue strength.

In the present invention, Al is prescribed as being 0.005% or less. This is because if over 0.005%, oxides of mainly Al₂O₃ would easily form, so fracture would result due to the oxides and sufficient fatigue strength and stability of quality would not be able to be secured. Further, when requiring a high fatigue strength, the content is preferably made 0.003% or less.

Further, one or more of Te, Sb, Mg, Zr, Ca, and Hf are added as elements for controlling the form of the oxides and sulfides when further higher performance and more stable performance are sought.

Te: 0.0002 to 0.01%

Te has the effect of making MnS spherical. If less than 0.0002%, the effect is not clear, while if over 0.01%, the matrix falls in toughness, hot cracking occurs, the fatigue durability is reduced, and other remarkable problems occur, so 0.01% is made the upper limit.

Sb: 0.0002 to 0.01%

Sb has the effect of making MnS spherical. If less than 0.0002%, the effect is not clear, while if over 0.01%, the matrix falls in toughness, hot cracking occurs, the fatigue durability is reduced, and other remarkable problems occur, so 0.01% is made the upper limit.

Mg: 0.0001 to 0.0005%

Mg forms oxides in molten steel of a temperature higher than the MnS formation temperature. These are already present in the molten steel at the time of MnS formation. Therefore, they can be used as nuclei for precipitation of MnS. Due to this, the distribution of MnS can be controlled. Further, looking at the number distribution as well, Mg-based oxides are dispersed in the molten steel more finely than the Si- and Al-based oxides often seen in conventional steel, so the MnS precipitated using the Mg-based oxides as nuclei finely disperses in the steel. Therefore, even with the same S content, the distribution of MnS differs depending on the presence or absence of Mg. Addition of these results in a finer MnS grain size. This effect is sufficiently obtained even in a small amount. If Mg is added, MnS is made finer. However, if exceeding 0.0005%, not only are hard oxides easily formed, but also MgS and other sulfides start to be formed so a drop in the fatigue strength and a drop in the coilability are invited. Therefore, the amount of addition of Mg was made 0.0001 to 0.0005%. When used for a high strength spring, an amount of 0.0003% or less is preferable. The amount of the element is small, but about 0.0001% can be added by making liberal use of Mg-based refractories. Further, Mg may be added by carefully selecting the secondary materials and using secondary materials with small Mg contents.

Zr: 0.0001 to 0.0005%

Zr is an oxide- and sulfide-forming element. Oxides are finely dispersed in the spring steel, so like Mg, form nuclei for precipitation of MnS. Due to this, the fatigue durability is improved and the ductility is increased to improve the coilability. If less than 0.0001%, this effect is not seen. Further, even if added over 0.0005%, formation of hard oxides is promoted, so even if the sulfides finely disperse, trouble due to oxides easily occurs. Further, with large addition, not only oxides, but also ZrN, ZrS, and other nitrides and sulfides are formed and cause trouble in production or a drop in the fatigue durability property of the spring, so the amount was made 0.0005% or less. Further, when using this for a high strength spring, the amount of addition is preferably made 0.0003% or less. These elements are small in amount, but can be controlled by careful selection of the secondary materials and precisely controlling the refractories etc.

For example, if making liberal use of Zr refractories in locations in contact with molten steel for a long time such as the ladle, tundish, nozzle, etc., it is possible to add 1 ppm or so with respect to 200 tons or so of molten steel. Further, while considering this, it is possible to add secondary materials so as not to exceed a prescribed range.

The method of analysis of Zr in the steel is to sample 2 g from the part of the steel material being measured free from the effect of surface scale, treat the sample by the same method as in Attachment 3 of JIS G 1237-1997, then measure it by ICP. At this time, the calibration line in ICP is set to be suitable for the fine amount of Zr.

Ca: 0.0002 to 0.01%

Ca is an oxide- and sulfide-forming element. In spring steel, by making MnS spherical, the length of MnS, which serves as a starting point of fatigue and other fracture, can be suppressed to make it harmless. The effect becomes unclear if less than 0.0002%. Further, even if added over 0.01%, not only is the yield poor, but also oxides or CaS or other sulfides are formed and cause trouble in production or a drop in the fatigue durability property of the spring, so the amount was made not more than 0.01%. The amount of addition is preferably not more than 0.001%.

Hf: 0.0002 to 0.01%

Hf is an oxide-forming element and forms the nuclei of precipitation of MnS. For this reason, Hf is an element forming oxides and sulfides by fine dispersion. In spring steel, the oxides are finely dispersed, so like Mg, form nuclei of precipitation of MnS. Due to this, the fatigue durability is improved and the ductility is increased to improve the coilability. This effect is not clear if the amount is less than 0.0002%. Further, even if over 0.01% is added, the yield is poor. Not only this, but also oxides or HfN, HfS, or other nitrides and sulfides are formed and cause production trouble or a drop in the fatigue durability property of the spring, so the amount was made 0.01% or less. This amount of addition is preferably 0.003% or less.

Below, the preferable ranges of content of other ingredients will be explained.

P: 0.015% or less

While P and S are not limited in the claims, they have to be restricted. P causes the steel to harden. Further, it segregates and makes the material brittle. The P segregating at the austenite grain boundaries causes a drop in the impact value and delayed fracture due to entry of hydrogen. For this reason, the smaller the amount the better. Therefore, P was limited to 0.015% or less where the embrittlement tends to become remarkable. Further, in the case of a high strength where the tensile strength of the heat treated steel wire exceeds 2150 MPa, a content of less than 0.01% is preferable.

S: 0.015% or less

S, like P, causes the steel to become brittle when present in steel. Mn reduces its effect sharply, but MnS also takes the form of inclusions, so lowers the fracture characteristic. In particular, in high strength steel, a small amount of MnS sometimes causes fracture. Therefore, it is preferable to reduce the S as much as possible. 0.015%, where this detrimental effect becomes remarkable, was therefore made the upper limit.

Further, in the case of a high strength where the tensile strength of the heat treated steel wire exceeds 2150 MPa, the amount is preferably made less than 0.01%.

t-O: 0.0002 to 0.01

In the steel, oxides formed through the deoxidation process and dissolved O are present. However, if the total amount of this oxygen (t-O) is large, it means that there are many oxide-based inclusions. If the oxide-based inclusions are small in size, the spring performance is not affected, but if a large amount of large oxides are present, the spring performance will be greatly affected.

If the amount of oxygen exceeds 0.01%, the spring performance is remarkably reduced, so the upper limit is preferably made 0.01%. Further, the smaller the amount of oxygen the better, but even if less than 0.0002%, the effect is saturated, so this is preferably made the lower limit.

If considering the ease of the actual deoxidation process etc., the content is preferably adjusted to 0.0005 to 0.005%.

Therefore, the following provision is added regarding the alloy-based spherical carbides and cementite-based spherical carbides on this examined surface. The following provisions are important for eliminating the harm due to these.

Prior austenite grain size number of #10 or more

In steel wire basically comprised of a tempered martensite structure, the prior austenite grain size has a major effect on the basic properties of the steel wire along with the carbides. That is, the smaller the prior austenite grain size, the better the fatigue characteristics and coilability. However, no matter how small the austenite grain size, if the carbides are included in a greater amount than prescribed, the effect is small. In general, to make the austenite grain size smaller, it is effective to reduce the heating temperature at the time of quenching, but this conversely increases the undissolved spherical carbides. Therefore, it is important to finish the steel wire into wire with a balance of the amount of carbides and prior austenite grain size. Here, when the carbides satisfy the above provision, if the prior austenite grain size number is less than #10, sufficient fatigue characteristics and coilability cannot be obtained, so the prior austenite grain size number is prescribed as being #10 or higher.

Further, for application to a high strength spring, a further finer grain size is preferable. By making it #11 or further #12 or more, both high strength and coilability can be achieved.

Residual Austenite of 15 Mass % or Less

Residual austenite often remains at the segregated parts, prior austenite grain boundaries, or near the regions sandwiched between sub grains. Residual austenite becomes martensite due to work-induced transformation. If induced transformation occurs at the time of spring fabrication, locally high hardness parts will be formed in the material and rather the coiling characteristics of the spring will be lowered. Further, recent springs are strengthened on their surfaces by shot peening, setting, and other plastic deformation, but in a production process including a number of such processes giving plastic deformation, the work-induced martensite formed at an early stage would lower the fracture strain and lower the workability and fracture characteristics of the spring during use. Further, even when industrially unavoidable deformation such as bruising is introduced, the material will easily fracture during coiling.

Further, by gradually breaking down in the nitridation, strain relief annealing, or other heat treatment, it will cause a change in the mechanical properties, lower the strength, reduce the coilability, and cause other trouble.

Therefore, the residual austenite is reduced as much as possible to suppress the formation of work-induced martensite and thereby improve the workability.

Specifically, if the amount of residual austenite exceeds 15% (mass %), the sensitivity to bruising etc. will become higher and the material will easily fracture during coiling or other handling, so the content was restricted to 15% or less.

The amount of residual austenite changes depending on the amounts of addition of C, Mn, and other alloy elements and the heat treatment conditions. For this reason, not only the design of the ingredients, but also the improvement of the heat treatment conditions are important.

When the martensite formation temperature (start temperature Ms point and finish temperature Mf point) becomes a low temperature, unless made a considerably low temperature at the time of quenching, martensite will not be formed and residual austenite will easily remain. In industrial quenching, water or oil is used, but to suppress residual austenite, sophisticated heat treatment control becomes necessary. Specifically, control becomes necessary to maintain the cooling medium at a low temperature, maintain an extremely low temperature even after cooling, secure a long transformation time to martensite, etc. Industrially, the material is processed on a continuous line, so the temperature of the cooling medium easily rises to close to 100° C., but it is preferable to maintain it at 60° C. or less. Further, a low temperature of 40° C. or less is more preferable. Further, to sufficiently promote martensite transformation, it is necessary to hold the material in the cooling medium for 1 second or more. Securing the holding time after cooling is also important.

Further, in addition to the provisions on the carbides etc., a structure in which the distribution of carbides is smaller than in other parts should be avoided. Specifically, in a lenticular martensite or its tempered structure, the distribution of carbides is smaller than in other parts and the microstructure is uneven, so the fatigue strength and workability are detrimentally affected.

EXAMPLES

Below, examples will be used to explain the present invention in detail.

To clarify the effects of the present invention, a material prepared in a 16 kg vacuum melting furnace was tested for its characteristics.

TABLE 1 (mass %) C Si Mn P S N V Cr Nb Mo W Ex. 0.61 2.37 0.57 0.004 0.009 0.0031 0.20 1.24 0.005 0.11 0.22

Steel of the chemical compositions shown in Table 1 was produced in a 16 kg vacuum melting furnace, forged, then normalized, patented, quenched, and tempered. The steel was normalized at two levels of temperature: (A) 1250° C.×1 hr→gradual cooling and (B) 1150° C.×1 hr→gradual cooling, then heated at 950° C. for patenting, then quenched and tempered at a quenching temperature of 910° C. and a tempering temperature of 500° C. After this, the materials of (A) and (B) were further annealed at 450, 475, and 500° C.×15 min. The tensile characteristics were evaluated -including for materials with no annealing. That is, a total of eight types of heat treated steel, including the materials of (A) and (B) with different normalization temperatures and with no annealing, were prepared and their tensile characteristics evaluated.

FIG. 2 shows the tensile characteristics of the tensile strength (shown by the solid line) and yield point (shown by the broken line) for +10 mm wire materials comprised of the material of Table 1 given the above heat treatments. Example (A) is a material normalized so that a large amount of V remains in the filtrate after filtering like in the present invention (soaking conditions 1250° C.×3 hr or more), while (B) is a material normalized so that undissolved carbides easily remain (soaking conditions 900° C.×30 min).

As a result, (A) and (B) remained almost the same in tensile strength and yield point as before annealing. Despite this, the [amount of V in filtrate filtered by 0.2 μm filter (mass %)] was 0.098% with (A) and 0.058% with (B). With (A), the annealing resulted in the tensile strength and yield point both becoming equal to or higher than those before annealing as a general trend, while with (B), the higher the annealing temperature, the lower both the tensile strength and yield point.

Both (A) and (B) remained almost the same in strength right after quenching and tempering. In this example, the yield point is relatively low and the coiling easy. However, depending on the nitridation after spring fabrication, with (A), the rise in the yield point was particularly remarkable. It was shown that when used as a spring, even if the hardness or tensile strength appeared low, the yield point was high, so the spring was superior in durability. That is, it was shown that the spring fabrication was easy and the subsequent nitridation raised the strength.

On the other hand, with (B), it was shown that the strength and yield point after spring fabrication fell along with the nitridation, so even if a nitride layer was formed at the surface layer, the inside was soft, local yield easily occurred, and the durability was inferior.

Preparation of Material and Heat History

Tables 2 to 7 show lists of examples and results of their evaluation. After that, the heat histories of the invention examples were shown. In part of the comparative examples, heat treatment different from that was performed, for example, treatment was performed under conventional general conditions deliberately so as to show the effects of the invention. These detailed conditions are explained in Tables 2 to 7.

TABLE 2 Chemical compositions (mass %) Ex. No. C Si Mn P S N V Cr Ti Nb Mo W Inv. ex. 1 0.70 2.08 0.66 0.006 0.006 0.0032 0.21 0.93 0.11 0.20 Inv. ex. 2 0.64 2.06 0.56 0.010 0.008 0.0038 0.21 0.84 0.12 0.11 Inv. ex. 3 0.52 2.31 0.79 0.006 0.006 0.0038 0.26 1.29 0.16 0.25 Inv. ex. 4 0.71 2.38 0.59 0.003 0.011 0.0034 0.21 1.25 0.008 0.18 0.19 Inv. ex. 5 0.63 2.00 0.67 0.005 0.007 0.0038 0.25 1.29 0.007 0.13 0.29 Inv. ex. 6 0.51 2.45 0.74 0.011 0.004 0.0037 0.22 0.93 0.009 0.14 0.18 Inv. ex. 7 0.46 2.25 0.65 0.009 0.005 0.0039 0.19 0.90 0.008 0.13 0.29 Inv. ex. 8 0.68 2.11 0.53 0.005 0.007 0.0039 0.24 0.99 0.008 0.15 0.34 Inv. ex. 9 0.70 2.37 0.75 0.009 0.007 0.0039 0.17 1.44 0.007 0.18 0.36 Inv. ex. 10 0.68 2.06 0.60 0.005 0.009 0.0040 0.12 1.36 0.008 0.14 0.19 Inv. ex. 11 0.70 2.00 0.64 0.009 0.005 0.0034 0.14 1.23 0.008 0.17 0.29 Inv. ex. 12 0.62 2.40 0.69 0.007 0.005 0.0033 0.45 0.75 0.008 0.02 0.15 Inv. ex. 13 0.61 1.65 0.67 0.011 0.010 0.0032 0.27 0.22 0.008 0.17 0.26 Inv. ex. 14 0.62 2.09 0.50 0.002 0.005 0.0032 0.25 0.22 0.008 0.13 0.19 Inv. ex. 15 0.70 2.13 0.64 0.008 0.008 0.0037 0.20 0.20 0.14 0.13 Inv. ex. 16 0.65 2.24 0.67 0.009 0.008 0.0031 0.23 0.29 0.17 0.29 Inv. ex. 17 0.52 2.48 0.51 0.005 0.007 0.0032 0.22 0.22 0.12 0.17 Inv. ex. 18 0.69 1.97 0.72 0.010 0.002 0.0035 0.22 0.25 0.008 0.14 0.34 Inv. ex. 19 0.60 1.95 0.55 0.010 0.007 0.0036 0.21 0.29 0.009 0.12 0.24 Inv. ex. 20 0.51 2.10 0.73 0.005 0.010 0.0033 0.18 0.90 0.009 0.11 0.15 Inv. ex. 21 0.46 2.50 0.56 0.009 0.009 0.0038 0.35 1.56 0.007 0.12 0.30 Inv. ex. 22 0.59 1.99 0.69 0.010 0.009 0.0037 0.29 1.23 0.008 0.11 0.11 Inv. ex. 23 0.73 2.05 0.79 0.005 0.005 0.0032 0.21 1.17 0.012 0.11 0.28 Inv. ex. 24 0.62 2.43 0.76 0.003 0.005 0.0039 0.23 1.08 0.008 0.16 0.10 Inv. ex. 25 0.50 2.02 0.50 0.008 0.011 0.0035 0.29 1.17 0.007 0.13 0.21 Inv. ex. 26 0.65 1.88 0.48 0.002 0.004 0.0036 0.32 Inv. ex. 27 0.69 2.46 0.49 0.003 0.010 0.0038 0.17 0.10 0.17 Inv. ex. 28 0.66 2.13 0.69 0.005 0.006 0.0036 0.25 0.73 0.005 0.11 0.24 Inv. ex. 29 0.67 1.97 0.58 0.005 0.005 0.0039 0.19 0.92 0.007 0.050 0.31 Inv. ex. 30 0.73 2.37 0.53 0.004 0.007 0.0039 0.20 0.58 0.008 0.15 Chemical compositions (mass %) Ex. No. Ta Ni Cu Co B Al Ca Zr Hf Te Sb Mg Inv. ex. 1 0.002 Inv. ex. 2 0.002 Inv. ex. 3 0.003 Inv. ex. 4 0.003 Inv. ex. 5 0.001 Inv. ex. 6 0.002 Inv. ex. 7 0.002 Inv. ex. 8 0.001 Inv. ex. 9 0.002 Inv. ex. 10 0.002 Inv. ex. 11 0.002 Inv. ex. 12 0.002 Inv. ex. 13 0.002 Inv. ex. 14 0.003 Inv. ex. 15 0.002 0.0004 Inv. ex. 16 0.002 0.0003 Inv. ex. 17 0.002 0.0002 Inv. ex. 18 0.002 0.0004 Inv. ex. 19 0.002 0.0002 Inv. ex. 20 0.003 0.0004 Inv. ex. 21 0.002 0.0005 Inv. ex. 22 0.002 0.0002 Inv. ex. 23 0.003 0.0003 Inv. ex. 24 0.002 0.0004 Inv. ex. 25 0.001 0.0004 Inv. ex. 26 0.001 Inv. ex. 27 0.002 Inv. ex. 28 0.001 Inv. ex. 29 0.002 Inv. ex. 30 0.002

TABLE 3 Quenching and tempering Processing conditions Tensile Radiant characteristics furnace Heating temperature Tensile Drawing Residual high Soaking Patenting Quenching Tempering strength ratio γ austenite Example No. frequency ° C. ° C. ° C. ° C. MPa % # % Inv. ex. 1 OT 1250 950 910 500 2199 50 12 7.9 Inv. ex. 2 OT 1250 950 910 490 2157 47 12 8.3 Inv. ex. 3 OT 1250 950 910 470 2199 45 10 7.1 Inv. ex. 4 OT 1250 950 910 500 2175 47 12 8.7 Inv. ex. 5 OT 1250 950 910 490 2180 47 12 7.0 Inv. ex. 6 OT 1250 950 910 470 2187 49 13 6.9 Inv. ex. 7 OT 1250 950 910 470 2158 45 12 6.8 Inv. ex. 8 OT 1250 950 910 500 2196 47 13 6.5 Inv. ex. 9 OT 1250 950 910 500 2191 48 13 6.3 Inv. ex. 10 OT 1250 950 910 490 2168 50 13 7.6 Inv. ex. 11 OT 1250 950 910 500 2163 49 12 7.1 Inv. ex. 12 OT 1250 950 910 480 2167 49 12 8.9 Inv. ex. 13 OT 1250 950 910 480 2199 46 13 7.2 Inv. ex. 14 OT 1250 950 910 490 2190 46 13 8.3 Inv. ex. 15 OT 1250 950 910 500 2151 49 11 7.8 Inv. ex. 16 OT 1250 950 910 480 2195 49 12 7.8 Inv. ex. 17 OT 1250 950 910 470 2166 47 11 7.4 Inv. ex. 18 OT 1250 950 910 490 2205 46 12 8.1 Inv. ex. 19 OT 1250 950 910 470 2162 45 13 6.7 Inv. ex. 20 OT 1250 950 910 470 2168 50 12 8.0 Inv. ex. 21 OT 1250 950 910 460 2171 47 13 6.2 Inv. ex. 22 OT 1250 950 910 470 2157 45 13 7.2 Inv. ex. 23 OT 1250 950 910 480 2165 48 12 8.8 Inv. ex. 24 OT 1250 950 910 480 2179 45 13 6.1 Inv. ex. 25 OT 1250 950 910 470 2159 50 12 6.6 Inv. ex. 26 OT 1250 950 910 490 2147 49 11 8.2 Inv. ex. 27 OT 1250 950 910 490 2170 46 11 8.3 Inv. ex. 28 OT 1250 950 910 480 2164 49 11 6.2 Inv. ex. 29 OT 1250 950 910 480 2183 47 14 6.3 Inv. ex. 30 OT 1250 950 910 510 2176 49 12 6.2 Annealing material Tensile Electrolyzed characteristics filtered V Amount Tensile Yield amount added of Hardness strength point Example No. % V × 0.4% HV MPa MPa Inv. ex. 1 0.107 0.083 538 1877 1785 Inv. ex. 2 0.109 0.085 521 1800 1744 Inv. ex. 3 0.137 0.103 564 1870 1804 Inv. ex. 4 0.109 0.085 585 1944 1851 Inv. ex. 5 0.130 0.101 557 1883 1817 Inv. ex. 6 0.111 0.086 554 1843 1777 Inv. ex. 7 0.106 0.078 536 1808 1728 Inv. ex. 8 0.127 0.094 560 1925 1867 Inv. ex. 9 0.088 0.067 583 1942 1871 Inv. ex. 10 0.054 0.012 546 1823 1766 Inv. ex. 11 0.076 0.058 545 1875 1782 Inv. ex. 12 0.242 0.180 549 1921 1847 Inv. ex. 13 0.145 0.109 507 1795 1722 Inv. ex. 14 0.133 0.099 506 1817 1749 Inv. ex. 15 0.108 0.080 528 1828 1749 Inv. ex. 16 0.127 0.093 537 1967 1889 Inv. ex. 17 0.117 0.088 520 1834 1767 Inv. ex. 18 0.118 0.089 536 1895 1836 Inv. ex. 19 0.107 0.082 516 1812 1742 Inv. ex. 20 0.092 0.070 527 1721 1651 Inv. ex. 21 0.181 0.140 575 1890 1805 Inv. ex. 22 0.155 0.114 543 1834 1767 Inv. ex. 23 0.111 0.084 556 1932 1853 Inv. ex. 24 0.122 0.090 571 1895 1827 Inv. ex. 25 0.155 0.114 531 1818 1742 Inv. ex. 26 0.166 0.128 484 1713 1658 Inv. ex. 27 0.086 0.067 527 1871 1783 Inv. ex. 28 0.137 0.102 548 1853 1779 Inv. ex. 29 0.101 0.075 542 1830 1741 Inv. ex. 30 0.106 0.081 544 1814 1745

TABLE 4 Chemical compositions (mass %) Example No. C Si Mn P S N V Cr Ti Nb Mo W Ta Inv. ex. 31 0.66 1.98 0.66 0.006 0.011 0.0036 0.22 0.75 0.008 0.32 Inv. ex. 32 0.55 2.07 0.47 0.010 0.004 0.0031 0.19 1.17 0.007 0.12 Inv. ex. 33 0.70 2.15 0.76 0.007 0.010 0.0038 0.26 0.98 0.008 0.13 Inv. ex. 34 0.60 2.05 0.78 0.004 0.009 0.0032 0.18 1.14 0.009 0.14 Inv. ex. 35 0.53 2.00 0.66 0.003 0.007 0.0036 0.20 0.75 0.007 0.10 Inv. ex. 36 0.72 2.34 0.49 0.008 0.007 0.0031 0.24 0.76 0.33 Inv. ex. 37 0.61 2.19 0.49 0.008 0.009 0.0033 0.27 1.04 0.30 Inv. ex. 38 0.51 2.00 0.58 0.002 0.005 0.0035 0.27 0.97 0.35 Inv. ex. 39 0.74 2.44 0.59 0.011 0.008 0.0033 0.22 0.75 0.12 Inv. ex. 40 0.61 1.86 0.70 0.010 0.011 0.0037 0.24 0.61 0.13 Inv. ex. 41 0.53 2.15 0.63 0.006 0.007 0.0034 0.19 1.29 0.12 Inv. ex. 42 0.54 1.99 0.63 0.004 0.010 0.0038 0.19 0.79 0.32 Inv. ex. 43 0.62 1.97 0.73 0.005 0.008 0.0038 0.18 0.63 0.11 0.15 0.05 Inv. ex. 44 0.62 1.98 0.68 0.006 0.003 0.0039 0.20 1.14 0.007 0.11 0.30 0.05 Inv. ex. 45 0.58 2.39 0.70 0.011 0.003 0.0034 0.20 0.61 0.14 Inv. ex. 46 0.63 2.48 0.52 0.006 0.008 0.0036 0.29 0.80 0.11 Inv. ex. 47 0.60 2.18 0.62 0.007 0.005 0.0038 0.20 0.89 0.013 0.17 Inv. ex. 48 0.64 2.39 0.70 0.009 0.009 0.0031 0.23 0.75 0.11 0.11 Inv. ex. 49 0.63 2.11 0.57 0.003 0.004 0.0039 0.18 0.96 0.15 0.13 Inv. ex. 50 0.59 2.18 0.58 0.006 0.008 0.0032 0.18 0.81 0.12 0.19 Inv. ex. 51 0.64 1.91 0.77 0.006 0.005 0.0032 0.26 1.17 0.10 0.16 Inv. ex. 52 0.61 2.19 0.58 0.003 0.009 0.0036 0.26 0.63 0.14 0.20 Inv. ex. 53 0.68 2.15 0.59 0.007 0.005 0.0033 0.23 0.75 0.14 0.32 Inv. ex. 54 0.63 1.99 0.50 0.007 0.011 0.0033 0.21 0.97 0.19 Inv. ex. 55 0.64 2.50 0.65 0.010 0.010 0.0038 0.26 0.90 0.16 0.33 Inv. ex. 56 0.60 2.01 0.69 0.011 0.003 0.0033 0.23 0.71 0.14 0.15 Inv. ex. 57 0.59 2.44 0.65 0.010 0.010 0.0031 0.17 1.22 0.17 0.15 Inv. ex. 58 0.74 2.37 0.34 0.011 0.005 0.0037 0.17 0.74 0.011 Inv. ex. 59 0.58 2.14 0.29 0.004 0.008 0.0032 0.28 0.65 0.007 0.17 0.15 Inv. ex. 60 0.60 2.18 0.28 0.004 0.008 0.0032 0.24 1.23 0.009 0.15 0.21 Chemical compositions (mass %) Example No. Ni Cu Co B Al Ca Zr Hf Te Sb Mg Inv. ex. 31 0.002 Inv. ex. 32 0.003 Inv. ex. 33 0.001 Inv. ex. 34 0.002 Inv. ex. 35 0.003 Inv. ex. 36 0.003 Inv. ex. 37 0.003 Inv. ex. 38 0.003 Inv. ex. 39 0.002 Inv. ex. 40 0.003 Inv. ex. 41 0.001 Inv. ex. 42 0.002 Inv. ex. 43 0.002 Inv. ex. 44 0.003 Inv. ex. 45 0.3 0.002 Inv. ex. 46 0.15 0.002 Inv. ex. 47 0.21 0.002 Inv. ex. 48 0.23 0.003 Inv. ex. 49 0.35 0.001 Inv. ex. 50 0.0018 0.003 Inv. ex. 51 0.002 Inv. ex. 52 0.002 0.005 Inv. ex. 53 0.002 0.0023 Inv. ex. 54 0.002 Inv. ex. 55 0.002 0.0005 Inv. ex. 56 0.003 0.012 Inv. ex. 57 0.002 0.0003 Inv. ex. 58 0.003 Inv. ex. 59 0.003 Inv. ex. 60 0.002

TABLE 5 Quenching and tempering Processing conditions Tensile Radiant characteristics furnace Heating temperature Tensile Drawing Residual high Soaking Patenting Quenching Tempering strength ratio γ austenite Example No. frequency ° C. ° C. ° C. ° C. MPa % # % Inv. ex. 31 OT 1250 950 910 490 2198 47 14 8.1 Inv. ex. 32 OT 1250 950 910 460 2169 50 13 7.0 Inv. ex. 33 OT 1250 950 910 500 2157 47 13 6.7 Inv. ex. 34 OT 1250 950 910 490 2154 47 14 7.3 Inv. ex. 35 OT 1250 950 910 470 2183 45 13 6.0 Inv. ex. 36 IQT 1250 950 950 500 2193 47 10 4.1 Inv. ex. 37 IQT 1250 950 950 480 2190 48 11 5.4 Inv. ex. 38 IQT 1250 950 950 470 2186 44 12 5.7 Inv. ex. 39 IQT 1250 950 950 510 2193 48 10 4.1 Inv. ex. 40 IQT 1250 950 950 480 2178 47 12 5.1 Inv. ex. 41 IQT 1250 950 950 470 2184 47 12 5.2 Inv. ex. 42 OT 1250 950 910 470 2150 48 10 8.4 Inv. ex. 43 OT 1250 950 910 480 2170 48 10 7.1 Inv. ex. 44 OT 1250 950 910 480 2178 50 13 7.5 Inv. ex. 45 OT 1250 950 910 470 2167 45 10 6.7 Inv. ex. 46 OT 1250 950 910 490 2151 50 10 6.4 Inv. ex. 47 OT 1250 950 910 480 2179 49 13 7.0 Inv. ex. 48 OT 1250 950 910 490 2183 47 11 6.6 Inv. ex. 49 OT 1250 950 910 480 2185 45 10 6.3 Inv. ex. 50 OT 1250 950 910 470 2156 47 11 6.4 Inv. ex. 51 OT 1250 950 910 480 2203 49 10 6.9 Inv. ex. 52 OT 1250 950 910 470 2164 50 11 7.7 Inv. ex. 53 OT 1250 950 910 490 2208 44 11 6.3 Inv. ex. 54 OT 1250 950 910 480 2177 47 10 7.4 Inv. ex. 55 OT 1250 950 910 490 2200 46 12 6.3 Inv. ex. 56 OT 1250 950 910 480 2189 45 10 8.7 Inv. ex. 57 OT 1250 950 910 480 2167 50 11 7.1 Inv. ex. 58 OT 1250 950 910 510 2153 50 11 8.7 Inv. ex. 59 OT 1250 950 910 470 2164 49 11 6.2 Inv. ex. 60 OT 1250 950 910 420 2329 45 12 6.9 Annealing material Tensile Electrolyzed characteristics filtered V Amount Tensile Yield amount added of Hardness strength point Example No. % V × 0.4% HV MPa MPa Inv. ex. 31 0.123 0.089 531 1815 1741 Inv. ex. 32 0.099 0.076 524 1687 1608 Inv. ex. 33 0.141 0.104 551 1879 1791 Inv. ex. 34 0.099 0.074 534 1744 1678 Inv. ex. 35 0.109 0.079 493 1693 1617 Inv. ex. 36 0.124 0.096 551 1861 1786 Inv. ex. 37 0.147 0.107 543 1778 1702 Inv. ex. 38 0.140 0.108 520 1758 1699 Inv. ex. 39 0.118 0.087 558 1908 1843 Inv. ex. 40 0.128 0.098 516 1723 1638 Inv. ex. 41 0.102 0.078 532 1707 1630 Inv. ex. 42 0.096 0.074 513 1869 1796 Inv. ex. 43 0.099 0.073 511 1808 1730 Inv. ex. 44 0.112 0.082 554 1835 1746 Inv. ex. 45 0.110 0.082 520 1835 1755 Inv. ex. 46 0.150 0.116 546 1880 1818 Inv. ex. 47 0.101 0.078 537 1823 1764 Inv. ex. 48 0.120 0.094 556 1889 1820 Inv. ex. 49 0.096 0.070 536 1843 1773 Inv. ex. 50 0.097 0.074 541 1851 1790 Inv. ex. 51 0.139 0.105 533 1806 1716 Inv. ex. 52 0.139 0.106 542 1895 1819 Inv. ex. 53 0.118 0.091 548 1910 1837 Inv. ex. 54 0.111 0.084 520 1715 1630 Inv. ex. 55 0.138 0.106 588 1955 1858 Inv. ex. 56 0.121 0.093 535 1792 1731 Inv. ex. 57 0.090 0.067 530 1840 1748 Inv. ex. 58 0.087 0.067 535 1769 170 Inv. ex. 59 0.145 0.112 534 1816 1744 Inv. ex. 60 0.122 0.094 522 1783 1720

TABLE 6 Chemical compositions (mass %) Example No. C Si Mn P S N V Cr Ti Nb Mo W Comp. ex. 61 0.70 2.17 0.71 0.004 0.011 0.0063 0.30 1.23 0.15 0.26 Comp. ex. 62 0.65 2.46 0.68 0.009 0.010 0.0064 0.30 1.29 0.17 0.21 Comp. ex. 63 0.51 2.18 0.75 0.007 0.009 0.0066 0.27 1.19 0.17 0.19 Comp. ex. 64 0.70 2.38 0.67 0.009 0.009 0.0036 0.24 1.02 0.009 0.14 0.21 Comp. ex. 65 0.57 2.03 0.63 0.006 0.006 0.0034 0.20 0.61 0.009 0.10 0.16 Comp. ex. 66 0.51 2.23 0.64 0.008 0.010 0.0040 0.20 1.24 0.008 0.16 0.20 Comp. ex. 67 0.67 2.08 0.65 0.010 0.011 0.0030 0.26 1.09 0.14 0.22 Comp. ex. 68 0.59 2.13 0.56 0.008 0.004 0.0033 0.20 0.98 0.11 0.12 Comp. ex. 69 0.61 2.23 0.65 0.008 0.004 0.0038 0.20 1.20 0.009 0.10 0.23 Comp. ex. 70 0.58 2.11 0.50 0.004 0.010 0.0034 0.29 0.84 0.16 0.28 Comp. ex. 71 0.62 2.36 0.49 0.007 0.003 0.0037 0.19 1.08 0.10 0.12 Comp. ex. 72 0.61 2.24 0.57 0.011 0.003 0.0033 0.26 0.75 0.009 0.15 0.18 Comp. ex. 73 0.58 2.42 0.74 0.003 0.003 0.0033 0.22 1.24 0.008 0.14 0.29 Comp. ex. 74 0.65 2.38 0.56 0.006 0.006 0.0036 0.19 1.10 0.13 0.19 Comp. ex. 75 0.64 2.12 0.73 0.010 0.010 0.0032 0.16 0.87 0.008 0.11 0.28 Comp. ex. 76 0.59 1.90 0.56 0.003 0.002 0.0034 0.20 1.22 0.009 0.13 0.17 Comp. ex. 77 0.57 2.44 0.76 0.009 0.004 0.0032 0.28 1.01 0.14 0.13 Comp. ex. 78 0.74 2.25 0.71 0.010 0.005 0.0034 0.26 0.63 0.16 Comp. ex. 79 0.65 2.38 0.61 0.005 0.004 0.0032 0.27 1.16 0.11 Comp. ex. 80 0.55 2.17 0.58 0.009 0.006 0.0033 0.18 0.87 0.12 Comp. ex. 81 0.62 2.06 0.72 0.010 0.005 0.0056 0.18 1.30 0.14 0.27 Comp. ex. 82 0.58 2.19 1.12 0.008 0.06 0.0035 0.27 0.88 0.12 0.24 Inv. ex. 83 0.61 2.27 0.57 0.005 0.005 0.0044 0.22 0.96 0.13 0.15 Inv. ex. 84 0.62 2.38 0.72 0.004 0.004 0.0038 0.20 1.21 0.007 0.13 0.23 Inv. ex. 85 0.72 2.14 0.84 0.003 0.003 0.0037 0.22 0.67 Chemical compositions (mass %) Example No. Ta Ni Cu Co B Al Ca Zr Hf Te Sb Mg Comp. ex. 61 0.003 Comp. ex. 62 0.002 Comp. ex. 63 0.001 Comp. ex. 64 0.003 Comp. ex. 65 0.001 Comp. ex. 66 0.002 Comp. ex. 67 0.002 Comp. ex. 68 0.002 Comp. ex. 69 0.001 Comp. ex. 70 0.003 Comp. ex. 71 0.002 Comp. ex. 72 0.002 Comp. ex. 73 0.003 Comp. ex. 74 0.001 Comp. ex. 75 0.002 Comp. ex. 76 0.003 Comp. ex. 77 0.002 Comp. ex. 78 0.003 Comp. ex. 79 0.003 Comp. ex. 80 0.001 Comp. ex. 81 0.001 Comp. ex. 82 0.002 Inv. ex. 83 0.02 0.003 Inv. ex. 84 0.02 0.003 Inv. ex. 85 0.014 0.003

TABLE 7 Quenching and tempering Processing conditions Tensile Radiant characteristics furnace Heating temperature Tensile Drawing Residual high Soaking Patenting Quenching Tempering strength ratio γ austenite Example No. frequency ° C. ° C. ° C. ° C. MPa % # % Comp. ex. 61 OT 1250 950 910 500 2173 36 10 8.7 Comp. ex. 62 OT 1250 950 910 470 2198 35 9 7.8 Comp. ex. 63 OT 1250 950 910 460 2199 34 10 6.2 Comp. ex. 64 OT 1250 890 910 490 2196 37 10 8.7 Comp. ex. 65 OT 1250 890 910 460 2150 38 10 7.4 Comp. ex. 66 OT 1250 890 910 460 2191 38 11 6.7 Comp. ex. 67 OT 1150 950 910 480 2202 38 12 7.6 Comp. ex. 68 OT 1150 950 910 460 2148 39 12 7.7 Comp. ex. 69 OT 1150 950 910 480 2177 39 11 7.7 Comp. ex. 70 OT 1150 950 910 460 2179 38 11 8.8 Comp. ex. 71 IQT 1250 950 900 480 2176 35 13 8.8 Comp. ex. 72 IQT 1250 950 900 380 2173 35 13 9.0 Comp. ex. 73 IQT 1250 950 900 470 2165 34 13 8.7 Comp. ex. 74 OT 1250 950 890 480 2200 33 13 7.4 Comp. ex. 75 OT 1250 950 890 470 2194 36 13 6.8 Comp. ex. 76 OT 1250 950 890 460 2158 37 13 7.4 Comp. ex. 77 OT 1250 950 890 460 2154 38 13 6.8 Comp. ex. 78 IQT 1250 950 900 510 2193 37 13 4.6 Comp. ex. 79 IQT 1250 950 900 490 2179 39 13 5.2 Comp. ex. 80 IQT 1250 950 900 460 2171 34 13 4.3 Comp. ex. 81 OT 1150 950 960 470 2176 37 8 7.8 Comp. ex. 82 OT 1250 950 910 460 2195 48 9 15.5 Inv. ex. 83 OT 1250 950 910 490 2152 47 10 10.1 Inv. ex. 84 OT 1250 950 910 480 2207 46 12 9.5 Inv. ex. 85 OT 1250 950 910 500 2153 50 11 8.7 Annealing material Tensile Electrolyzed characteristics filtered V Added Tensile Yield amount amount of Hardness strength point Example No. % V × 0.4% HV MPa MPa Comp. ex. 61 0.095 0.118 575 1882 1719 Comp. ex. 62 0.083 0.119 585 1932 1763 Comp. ex. 63 0.090 0.106 548 1768 1607 Comp. ex. 64 0.076 0.096 581 1895 1624 Comp. ex. 65 0.068 0.082 507 1705 1530 Comp. ex. 66 0.051 0.079 553 1701 1477 Comp. ex. 67 0.068 0.105 558 1782 1629 Comp. ex. 68 0.061 0.078 538 1706 1564 Comp. ex. 69 0.069 0.082 552 1770 1527 Comp. ex. 70 0.087 0.115 553 1752 1608 Comp. ex. 71 0.050 0.077 538 1782 1541 Comp. ex. 72 0.083 0.106 542 1771 1546 Comp. ex. 73 0.064 0.089 562 1817 1546 Comp. ex. 74 0.050 0.074 558 1826 1664 Comp. ex. 75 0.060 0.064 552 1798 1558 Comp. ex. 76 0.069 0.078 543 1662 1494 Comp. ex. 77 0.087 0.111 563 1794 1582 Comp. ex. 78 0.077 0.104 545 1860 1647 Comp. ex. 79 0.086 0.109 559 1818 1614 Comp. ex. 80 0.057 0.072 515 1653 1491 Comp. ex. 81 0.095 0.073 563 1738 1573 Comp. ex. 82 0.169 0.108 502 1624 1476 Inv. ex. 83 0.135 0.087 521 1838 1711 Inv. ex. 84 0.121 0.082 548 1840 1736 Inv. ex. 85 0.087 0.067 535 1769 1700

The materials of the present invention were produced by (a) a 270 ton converter (Example 5), (b) a 2 ton vacuum melting furnace (Examples 19, 39, and 59), and (c) a 16 kg vacuum melting furnace (other examples). The materials produced by the 270 ton converter and the 2 ton vacuum melting furnace were heated to 1250 to 1300° C. and rolled to prepare billets. At this time, the temperature was sufficiently raised so as to make the structures of the billets uniform and to make the V and other carbide-forming elements sufficiently dissolve.

Further, the billets were rolled to prepare the spring-use steel wire materials. At this time, in the invention examples, the materials were held for a constant time at high temperatures of 1200° C. or more. After this, in each case, the billets were rolled to φ8 mm.

In the other examples, the materials were melted by a 16 kg vacuum melting furnace, then forged to φ13 mm×600 mm, then heat treated. After this, they were held at 1300° C.×3 hr or more to again make the V and other carbide-forming elements sufficiently dissolve. After this, the materials were again held at high temperatures of 1200° C. or more for a certain time.

In the cooling process at the time of rolling or after rolling simulation, to suppress the formation of martensite or other hard, easily crackable supercooled structures, the material was cooled from the high temperature at the time of rolling to 450° C., then was gradually cooled by covering it etc. Even if a supercooled structure resulted from this, softening was possible, no cracks or defects were caused in the later processes as well, and handling was easy.

Heat Treatment of Wire Materials (Pretreatment)

Materials rolled or forged and given a heat history in the above way were patented, drawn, quenched, and tempered.

The materials were patented by heating at a temperature of 950° C.×20 min, then charged into a 600° C. Pb tank and converted to a ferrite pearlite structure. At this time, the pearlite transformation was ended in the patenting tank in an extremely short time.

In this state, the materials were drawn by dies to +4 mm.

Quenching and Tempering

The quenching and tempering were performed by both (1) radiant furnace heating and (2) high frequency induction heating.

(1) Radiant Furnace Heating

A material is charged into a radiant furnace heated to 920° C., then 10 minutes later is charged into a 50° C. oil tank and quenched. After 5 minutes, it is pulled up and in that state is charged into a Pb tank adjusted to a predetermined temperature and tempered. The temperature of a lead tank is 450 to 550° C., but in the invention examples was about 500° C. This was to avoid undissolved carbides and cause the precipitation of a large amount of extremely fine V-based carbides and was to lengthen as much as possible the holding time in the V precipitation temperature of 450 to 650° C.

(2) High Frequency induction Heating

In high frequency induction heating, a material is placed in a coil and heated to 900 to 1000° C. After induction heating, it is immediately water cooled. It can then be again heated in the coil to 400 to 600° C. to temper it.

To obtain the same strength, high frequency tempering enables treatment at a high temperature. After tempering, the material is water cooled.

After quenching and tempering, part is measured for the drawing ratio, an indicator of tensile strength and workability, by a tensile test.

At this time, the tempering temperature is controlled to give a tensile strength over 2100 MPa.

Nitridation Annealing

After quenching and tempering, each material was annealed assuming nitridation. When not particularly shown, it was held at 500° C.×2 hr. After this, the amount of V was measured using a filter. Due to the annealing assuming nitridation, the V-based carbides become larger, so a greater amount of V is detected on the filter compared with before annealing. Therefore, when the amount of V in filtrate filtered by a 0.2 μm filter (mass %) detected before annealing is greater than that after quenching and tempering and the provision of the present invention is satisfied by measurement of the amount of V in filtrate filtered by a 0.2 μm filter (mass %) after annealing, the heat treated steel before annealing also satisfies the provision of the present invention.

Evaluated Matter

The matter evaluated were as follows:

(1) After quenching and tempering: tensile strength, drawing ratio (workability), prior austenite grain size, amount of residual austenite

(2) After annealing: [amount of V in filtrate filtered by 0.2 μm filter (mass %)], hardness, tensile strength, yield point

The tensile test after the quenching and tempering measured the drawing ratio—an indicator of tensile strength and spring workability. Each material was quenched and tempered to exceed 2100 MPa, then a JIS Z 2201 No. 9 test piece was prepared and was tested in accordance with JIS Z 2241. The tensile strength was calculated from the fracture load.

Further, in recent years, springs have often been made higher in strength by hardening by nitridation of the surface layer. Nitridation heats a spring in a nitridation atmosphere gas to 400 to 500° C. and holds it there for several minutes to 1 hour or so in order to make the surface layer harden. At that time, the inside where the nitrogen has not penetrated is heated, so is annealed and softened. Suppressing this softening is important. The hardness (indicator of softening resistance), tensile strength, and yield point after annealing simulating nitridation were evaluated.

Nitrided and annealed steel is a material similar to the inside of a spring. Its high yield point means superior spring durability. Further, in actual springs, shot peening to impart compressive residual stress is becoming a general practice, but the compressive residual stress becomes greater proportionally to the yield point. A large yield point means a large compressive residual stress and further a deep residual stress layer as well. In this way, the ease by which compressive residual stress can remain is also one factor increasing the durability of an actual spring.

Regarding the [amount of V in filtrate filtered by 0.2 μm filter (mass %)], as already explained, this was calculated by measuring the amount of V on the filter after electrolysis and filtration and subtracting it from the amount of V added.

Tensile Test Method

The tensile test was performed in accordance with the JIS. A strain meter was attached and pulled to measure both the yield point and tensile strength. When the yield point was unclear, the 0.2% proof strength was measured as the yield point. Further, the drawing ratio was measured and used as an indicator for evaluating the workability.

Explanation Regarding Yield Point

Up until now, it had been believed that raising the strength meant raising the spring durability, but in fact the initial growth of a fracture is dependent on the yield of the material, so is dependent not only on the tensile strength in a tensile test, but also the yield point. Due to this, as for example shown in Uemura, Ominami, Introduction to Dynamics of Fracture, Ohm, Dec. 10, 1984, first edition, sixth printing, p. 102, even if the tensile strength is low, when the yield point is high, a high spring durability can be obtained.

Further, in recent spring steel, shot peening is often used to impart compressive residual stress to the surfacemost layer. The residual stress is caused by the unevenness of the high strain parts at the surface layer with the low strain parts at the inside. Therefore, regarding the size of the residual stress, the size of the yield point may be considered the maximum value of the actual residual stress. Therefore, the steel wire of the present invention, which can be raised in yield point by annealing, can be adjusted to a material suitable for shot peening by heat treatment.

Explanation of Results

Summarizing the results of Tables 2 to 7, the invention examples have chemical ingredients in the ranges shown in the claims. By keeping N low, the amount of V added can be stably increased. As a result, as shown by the measurement value of [amount of V in filtrate filtered by 0.2 μm filter (mass %)], the V-based undissolved carbides can be suppressed. Therefore, there is an effect on the tensile strength and yield point after annealing. Compared with the comparative examples, the tensile strength tends to be somewhat higher. Further, even if the tensile strength is equal, the yield point tends to be higher in the invention examples.

This difference is caused due in part to the effect of the N and other chemical ingredients such as in Examples 61 to 63, but the effect of the heat treatment is also large. That is, in production of steel, after casting, V-based carbides precipitate at the time of cooling. However, at the time of heating, if not heating to a temperature higher than the precipitation temperature at the time of cooling, the once precipitated precipitates will not dissolve. For this reason, undissolved carbides remain. To avoid this, it is important to maintain it at an easily dissolvable state at the heating temperature of the next process. Therefore, in the heating such of the rolling etc. after casting, treatment at an extremely high temperature is necessary.

In the present invention, therefore, the heating temperature in the rolling at the time of making the billet (blooming), the rolling into a rod product (wire rolling), the patenting for drawing, and the quenching and tempering to impart strength is maintained extremely high to promote the dissolution of the undissolved carbides in the different processes and suppress their growth. Specifically, as shown in the tables, the material was heated at a high temperature in each process.

Therefore, it is possible to secure strength after quenching and tempering of course and the drawing ratio became a value of no problem in spring fabrication. With φ4 mm heat treated steel wire, coiling is possible if the drawing ratio is over 45%.

Further, the [amount of V in filtrate filtered by 0.2 μm filter (mass %)] became 50% or so of the amount of V added. Precipitation is possible at the time of annealing. In Tables 2 to 7, the values are at the time of measurement after annealing, so the sum of the amount of V dissolved and the amount of V in the fine precipitates is measured. The annealing results in the formation of V-based precipitates, so before annealing a further greater amount of dissolved V is present, so a further greater amount of V is detected.

Examples 64 to 80 are examples of low heating temperatures for the blooming, patenting, and imparting of strength and where the V-based carbides could not be sufficiently dissolved.

Radiant furnace heating and high frequency induction heating differ in the holding time at the time of heating, so the heating temperatures differ. High frequency induction heating tends to heat to a higher temperature. With a radiant furnace, the temperature is 910° C. or so, while with high frequency induction heating, it is necessary to heat to a temperature of 950° C. or so, higher than with radiant furnace heating, in a short time for dissolution. On the other hand, in the tempering process, the heating can be ended in a short time, so the tempering can also be completed by a higher temperature in a shorter time than with radiant furnace heating.

When giving a different heat history in this way, in the heating up to the rolling and quenching, unless extremely high temperature heating is maintained, the undissolved carbides will grow, so will not be able to be sufficiently dissolved in the next process. As a result, the drawing ratio after quenching and tempering is inferior to the invention examples. This shows that working into a spring would be difficult. Further, after annealing, the [amount of V in filtrate filtered by 0.2 μm filter (mass %)] is smaller than in the invention examples, so the tensile strength and yield point after annealing tend to be low. Since the yield point is low, the spring is insufficient in durability. A durability of the level of the invention examples cannot be obtained.

Further, as in Example 81, even if the [amount of V in filtrate filtered by 0.2 μm filter (mass %)] is in the prescribed range, if forcibly heating at a high temperature to cause dissolution in the heating at the time of quenching, the prior austenite grain size will become larger, the drawing ratio after quenching and tempering and the yield point after annealing cannot be sufficiently secured, and both the workability and spring performance are inferior to those of the invention examples.

A similar trend can be seen in Example 82 where even if a large amount of residual austenite remains, the softening due to the annealing is extensive and a sufficient yield point can no longer be obtained.

INDUSTRIAL APPLICABILITY

The present invention steel not only reduces the austenite grain size and amount of residual austenite, but also suppresses the prescribed fine V-based carbides of 0.2 μm or less, which had a high possibility of being overlooked in the past, so as to enable steel wire increased in strength to a strength of 2000 MPa or more to exhibit better spring performance by nitridation or other heat treatment after spring fabrication. Due to this, it is possible to produce a spring securing coilability, high in strength and superior in fracture characteristics, and exhibiting a higher performance by nitridation or other heating. For this reason, the present invention steel has the remarkable effect of giving excellent properties to a spring even after being worked into a spring and enabling the characteristics of the material to be exhibited to the fullest. 

1. Spring steel characterized by containing, by mass %, C, 0.45 to 0.9%, Si: 1.0 to 3.0%, Mn: 0.1 to 2.0%, V: over 0.1 to 1.0%, N: limited to 0.007% or less, and a balance of Fe and unavoidable impurities, and satisfying, in terms of the analyzed value of the extracted residue after heat treatment, [amount of V in filtrate filtered by 0.2 μm filter(mass %)]≧[amount of V in steel(mass %)]×0.4.
 2. Spring-use heat treated steel wire as set forth in claim 1 characterized by further containing, by mass %, one or more of Cr: 0.5 to 2.5%, Nb: 0.001 to less than 0.05%, Ti: 0.001 to less than 0.05%, W: 0.05 to 0.5%, Mo: 0.05 to 0.5%, Ta: 0.001 to 0.5%, Ni: 0.05 to 3.0%, Cu: 0.05 to 0.5%, Co: 0.05 to 3.0%, B: 0.0005 to 0.006%, Te: 0.0002 to 0.01%, Sb: 0.0002 to 0.01%, Mg: 0.0001 to 0.0005%, Zr: 0.0001 to 0.0005%, Ca: 0.0002 to 0.01%, and Hf: 0.0002 to 0.01%.
 3. Spring-use heat treated steel wire as set forth in claim 1 characterized further by restricting, by mass %, Al to 0.005% or less.
 4. Heat treated steel for a high strength spring said characterized by having the steel ingredients as set forth in any one of claim 1, having a prior austenite grain size number after quenching and tempering of #10 or more, and having residual austenite in an amount of 15 mass % or less. 